A nickel-based alloy

ABSTRACT

A nickel-based alloy composition consisting, in weight percent, of: between 9.0 and 13.2% chromium, between 5.9 and 24.9% cobalt, between 0.0 and 4.0% iron, between 1.1 and 4.4% molybdenum, between 0.0 and 8.0% tungsten, between 2.8 and 3.7% aluminium, between 0.3 and 5.1% titanium, between 0.0 and 4.0% niobium, more than 2.4% tantalum and 9.5% or less tantalum, between 0.01 and 0.1% carbon, between 0.001 and 0.1% boron, between 0.001 and 0.3% zirconium, between 0.0 and 0.5% silicon, between 0.0 and 0.1% yttrium, between 0.0 and 0.1% lanthanum, between 0.0 and 0.1% cerium, between 0.0 and 0.003% sulphur, between 0.0 and 0.25% manganese, between 0.0 and 0.5% vanadium, between 0.0 and 0.5% copper, and between 0.0 and 0.5% hafnium, the balance being nickel and incidental impurities, wherein the following equations are satisfied in which WMO, Ww, WNb, W-Ta and WTi are the weight percent of molybdenum, tungsten, niobium, tantalum and titanium in the alloy respectively 1.29WMo+0.5Ww≥5.7, 0.6WTi+0.44WNb+0.27 WTa≥4.2.

The present invention relates to a nickel-based superalloy composition for use as a turbine disc component within a gas turbine engine and other turbomachinery. The turbine disc is a critical component in gas turbine engines. Increases in turbine disc alloy performance—in terms of maximum operating temperature and maximum service—life can have a significant impact on the efficiency of the engine as well as the cost effectiveness of operating the engine.

Examples of typical compositions of nickel-based superalloys which are used for turbine disc components are listed in Table 1. In development of higher strength alloys there has been a tendency to move towards higher levels of cost, FIG. 1.

TABLE 1 Nominal composition in wt. % of commonly applied nickel-based superalloys used for powder metallurgy turbine discs. Alloy (wt. %) Cr Co Fe Mo W Al Ti Ta Nb Hf C B Zr N18 11.50 15.70 0.00 6.50 0.00 4.35 4.35 0.00 0.00 0.50 0.015 0.015 0.030 René88DT 16.00 13.00 0.00 4.00 4.00 2.10 3.70 0.00 0.70 0.00 0.030 0.015 0.030 RR1000 15.00 18.50 0.00 5.00 0.00 3.00 3.60 2.00 0.00 0.50 0.030 0.020 0.060 ME3 13.00 20.50 0.00 3.70 2.00 3.40 3.60 2.40 0.90 0.00 0.040 0.030 0.050 Alloy10 10.46 17.93 0.00 2.52 4.74 3.53 3.79 1.61 0.97 0.00 0.027 0.028 0.070 LSHR 13.00 20.50 0.00 2.75 4.55 3.50 3.50 1.70 1.55 0.00 0.030 0.030 0.050 U.S. Pat. No. 8,147,749 10.09 19.60 0.00 2.79 2.62 3.14 2.17 7.28 1.55 0.40 0.030 0.030 0.050 N19 12.90 11.80 0.00 4.70 3.20 2.50 3.80 0.00 1.60 0.30 0.022 0.015 0.060 U.S. Pat. No. 8,613,810 12.00 18.00 0.00 2.90 2.80 3.20 3.10 5.10 1.50 0.40 0.055 0.025 0.055 US2015/0192002 13.40 14.00 0.00 1.30 5.30 3.25 3.10 3.70 1.55 0.00 0.050 0.025 0.055

It is the aim of the present invention to deliver high high-temperature strength and creep resistance in combination with a reasonable, preferably a reduction in, alloy cost and an improvement in oxidation/corrosion resistance. The balance of properties for the new alloy may make it more cost effective for the production of components for high temperature applications; in particular for use in a turbine disc applications where the operating temperature for the component is 800° C. or greater.

The present invention provides a nickel-based alloy composition consisting, in weight percent, of: between 9.0 and 13.2% chromium, between 5.9 and 24.9% cobalt, between 0.0 and 4.0% iron, between 1.1 and 4.4% molybdenum, between 0.0 and 8.0% tungsten, between 2.8 and 3.7% aluminium, between 0.3 and 5.1% titanium, between 0.0 and 4.0% niobium, more than 2.4% tantalum and 9.5% or less tantalum, between 0.01 and 0.1% carbon, between 0.001 and 0.1% boron, between 0.001 and 0.3% zirconium, between 0.0 and 0.5% silicon, between 0.0 and 0.1% yttrium, between 0.0 and 0.1% lanthanum, between 0.0 and 0.1% cerium, between 0.0 and 0.003% sulphur, between 0.0 and 0.25% manganese, between 0.0 and 0.5% vanadium, between 0.0 and 0.5% copper, and between 0.0 and 0.5% hafnium, the balance being nickel and incidental impurities, wherein the following equations are satisfied in which W_(Mo), W_(w), W_(Nb), W_(Ta) and W_(Ti) are the weight percent of molybdenum, tungsten, niobium, tantalum and titanium in the alloy respectively

1.29W _(Mo)+0.5W _(w)≥5.7

0.6W _(Ti)+0.44W _(Nb)+0.27W _(Ta)≥4.2.

This alloy provides exceptionally high strength at elevated temperatures, combined with excellent oxidation/corrosion resistance and resistance to TCP formation.

In a preferred embodiment the following equation is satisfied in which W_(Nb), W_(Ta), W_(Ti) and W_(Al) are the weight percent of niobium, tantalum, titanium and aluminium in the alloy respectively

6.2≤0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta) +0.94W _(Al)≤7.4

preferbly 6.2≤0.6W _(Ti)0.31W _(Nb)+0.15W _(Ta)+0.94W _(Al)≤6.7

This ensures the alloy has enough gamma prime phase to give a high strength.

In a preferred embodiment the following equation is satisfied in which W_(Nb), W_(Ta), W_(Ti) and W_(Al) are the weight percent of niobium, tantalum, titanium and aluminium in the alloy respectively

1.0≤(0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta))/W_(Al)23 1.3

This ensures an alloy with a high APB energy and stable gamma prime phase such that the alloy has good resistance to precipitate shearing and resistance to formation of eta and delta phases.

In a preferred embodiment the following equation is satisfied in which W_(W) and W_(Mo) are the weight percent of tungsten and molybdenum, in the alloy respectively

6.0≤1.29W _(Mo)+0.5W _(W)

Such an alloy has a high solid solution merit index meaning improved resistance to high temperature creep and improved high temperature strength.

In a preferred embodiment the following equation is satisfied in which W_(Ta) and W_(W) are the weight percent of tantalum and tungsten, in the alloy respectively

W _(Ta)+0.88W _(W)≤9.4

Such an alloy achieves a lower density than if the equation is not satisfied.

In a preferred embodiment the following equation is satisfied in which W_(Ta) and W_(Co) are the weight percent of tantalum and cobalt, in the alloy respectively

W _(Ta)+0.45W _(Co)≤14.1

The cost of such an alloy is limited.

In a preferred embodiment the following equation is satisfied in which W_(W), W_(Mo) and W_(Cr) are the weight percent of tungsten and chromium and molybdenum, in the alloy respectively

W _(W)+1.62W _(Cr)+2.0W _(Mo)≤30.2

preferably W _(W)+1.62W _(Cr)+2.0W _(Mo)≤28.3

Such an alloy has high resistance to the formation of deleterious TCP phases.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 3.5 wt. % or less aluminium, preferably 3.1 wt. % or less aluminium. Such an alloy has improved manufacturability due to a lower gamma prime solvus temperature.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 2.45 wt. % or more tantalum, preferably 2.5 wt. % or more tantalum, more preferably 3.4 wt. % or more tantalum, even more preferably 5.1 wt. % or more tantalum. Such an alloy has even higher strength at the cost of lower gamma prime solvus temperature.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 6.4 wt. % or less tantalum. Such an alloy maintains easy manufacturability due to limited gamma prime solvus temperature.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 1.0 wt. % or more niobium. Such an alloy has increased strength.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 3.0 wt. % or less niobium, preferably 2.0 wt. % or less niobium. Such an alloy has improved strength.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 1.7 wt. % or more of titanium, preferably 2.5 wt. % or more titanium. Such an alloy has increased strength.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 4.8 wt. % or less titanium, preferably 4.6 wt. % or less titanium, more preferably 4.2 wt. % or less titanium. Such an alloy has better oxidation resistance and improved manufacturability.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 3.2 wt. % or more tungsten. Such an alloy has excellent high temperature strength.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 7.8 wt. % or less tungsten, preferably 6.9 wt. % or less tungsten. Such an alloy has a lower density.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 1.2 wt. % or more molybdenum, preferably 1.7 wt. % or more molybdenum. Such an alloy has higher solid solution strengthening meaning improved creep strength.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 6.1 wt. % or more cobalt, preferably 7.1 wt. % or more cobalt. Such an alloy has increased strength, particularly creep resistance.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 24.6 wt. % or less cobalt, preferably 22.8 wt. % or less cobalt, more preferably 19.1 wt. % or even 16.8wt % or 16.6 wt. % or less cobalt, even more preferably 14.7 wt. % or less cobalt, most preferably 11.1 wt. % or less cobalt. Such an alloy has a lower cost.

In a preferred embodiment the nickel-based alloy composition consists of, in weight percent, 11.9 wt. % or less chromium, preferably 11.5 wt. % or less chromium, more preferably 11.2 wt. % or less chromium. Such an alloy has improved microstructural stability in terms of resistance to the generation of TCP phases.

The term “consisting of” is used herein to indicate that 100% of the composition is being referred to and the presence of additional components is excluded so that percentages add up to 100%. Unless otherwise stated, percent's are expressed in weight percent.

The invention will be more fully described, by way of example only, with reference to the accompanying drawings in which:

FIG. 1 shows the calculated trade-off between the elemental cost and yield strength (in terms of strength merit index) for the alloys listed in Table 1 and alloys within the alloy design space listed in Table 2, a hatched area on the graph designates a preferred combination of strength and cost;

FIG. 2 shows the calculated correlation between yield strength (in terms of strength merit index) and volume fraction of γ′ phase at 850° C. The minimum volume fraction of γ′ required to achieve the desired level of strength merit index is highlighted;

FIG. 3 is a contour plot showing the effect of elements aluminium and the sum of elements niobium, titanium and tantalum (according to the relationship 0.6W_(Ti)+0.31W_(Nb)+0.15W_(Ta)) on volume fraction of γ′ at 900° C. Also delinated on the graph are composition dependent relationships for γ′ solvus and strengthening;

FIG. 4 is a contour plot showing the effect of elements aluminium and the sum of elements niobium, titanium and tantalum (according to the relationship 0.6 W_(Ti)+0.31W_(Nb)+0.15 W_(Ta)) on γ′ solvus temperature;

FIG. 5 is a contour plot showing the effect of titanium and tantalum on strength merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the aluminium content is fixed between 2.8-3.7wt. % and niobium content is fixed at 0.0 wt. %;

FIG. 6 is a contour plot showing the effect of titanium and tantalum on strength merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the aluminium content is fixed between 2.8-3.7wt. % and niobium content is fixed at 1.0 wt. %;

FIG. 7 is a contour plot showing the effect of titanium and tantalum on strength merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the aluminium content is fixed between 2.8-3.7wt. % and niobium content is fixed at 2.0 wt. %;

FIG. 8 is a contour plot showing the effect of titanium and tantalum on strength merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the aluminium content is fixed between 2.8-3.7wt. % and niobium content is fixed at 3.0 wt. %;

FIG. 9 is a contour plot showing the effect of titanium and tantalum on strength merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the aluminium content is fixed between 2.8-3.7wt. % and niobium content is fixed at 4.0 wt. %;

FIG. 10 is a contour plot showing the effect of molybdenum and tungsten on solid solution merit index, for alloys with a volume fraction of γ′ between 0.51 and 0.62;

FIG. 11 is a contour plot showing the effect of tantalum, and tungsten on alloy density, for alloys with a volume fraction of γ′ between 0.51 and 0.62;

FIG. 12 is a contour plot showing the effect of cobalt and tantalum on alloy cost (in terms of raw material cost, based on 2017 elemental prices), for alloys with a volume fraction of γ′ between 0.51 and 0.62;

FIG. 13 is a contour plot showing the effect of elements cobalt and tungsten on creep resistance (in terms of creep merit index), for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 1.0 wt. %;

FIG. 14 is a contour plot showing the effect of elements cobalt and tungsten on creep resistance (in terms of creep merit index), for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 2.0 wt. %;

FIG. 15 is a contour plot showing the effect of elements cobalt and tungsten on creep resistance (in terms of creep merit index), for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 3.0 wt. %;

FIG. 16 is a contour plot showing the effect of elements cobalt and tungsten on creep resistance (in terms of creep merit index), for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 4.0 wt. %;

FIG. 17 is a contour plot showing the effect of elements cobalt and tungsten on creep resistance (in terms of creep merit index), for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 5.0 wt. %;

FIG. 18 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 0.0 wt. %;

FIG. 19 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 1.0 wt. %;

FIG. 20 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 2.0 wt. %;

FIG. 21 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 3.0 wt. %;

FIG. 22 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 4.0 wt. %; and

FIG. 23 is a contour plot showing the effect of elements chromium and tungsten on the stability number Md, for alloys with a volume fraction of γ′ between 0.51 and 0.62 when the molybdenum content is fixed at 5.0 wt. %.

Traditionally, nickel-based superalloys have been designed through empiricism. Thus their chemical compositions have been isolated using time consuming and expensive experimental development, involving small-scale processing of limited quantities of material and subsequent characterisation of their behaviour. The alloy composition adopted is then the one found to display the best, or most desirable, combination of properties. The large number of possible alloying elements indicates that these alloys are not entirely optimised and that improved alloys are likely to exist.

In superalloys, generally additions of chromium (Cr) and aluminium (Al) are added to impart resistance to oxidation/corrosion, cobalt (Co) is added to improve resistance to sulphidisation. For creep resistance, molybdenum (Mo), tungsten (W) and cobalt (Co) are introduced, because these retard the thermally-activated processes—such as, dislocation climb —which determine the rate of creep deformation. To promote static and cyclic strength, aluminium (Al), tantalum (Ta), niobium (Nb) and titanium (Ti) are introduced as these promote the formation of the precipitate hardening phase gamma-prime (γ′). This precipitate phase is coherent with the face-centered cubic (FCC) matrix phase which is referred to as gamma (γ).

A modelling-based approach used for the isolation of new grades of nickel-based superalloys is described here, termed the “Alloys-By-Design” (ABD) method. This approach utilises a framework of computational materials models to estimate design relevant properties across a very broad compositional space. In principle, this alloy design tool allows the so called inverse problem to be solved; identifying optimum alloy compositions that best satisfy a specified set of design constraints.

The first step in the design process is the definition of an elemental list along with the associated upper and lower compositional limits. The compositional limits for each of the elemental additions considered in this invention—referred to as the “alloy design space”—are detailed in Table 2.

TABLE 2 Alloy design space studied. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Min 6.00 0.00 0.00 0.00 2.00 0.00 0.00 0.00 0.050 0.025 0.055 Max 18.00 30.00 8.00 12.00 5.00 8.00 12.00 4.00

The second step relies upon thermodynamic calculations used to calculate the phase diagram and thermodynamic properties for a specific alloy composition. Often this is referred to as the CALPHAD method (CALculate PHAse Diagram). These calculations are conducted at the typical service temperature for the new alloy (900° C.), providing information about the phase equilibrium (microstructure).

A third stage involves isolating alloy compositions which have the desired microstructural architecture. In the case of nickel based superalloys which require superior resistance to creep deformation, the creep rupture life generally improves as the volume fraction of the precipitate hardening phase γ′ is increased, the most beneficial range for volume fraction of γ′ lies between 60%-70% at 900° C. (however often due to other design restraints volume fraction may be limited to lower values than this and so alloys with a γ′ volume fraction of 50% to 60% are included). At values above 70% volume fraction of γ′ a drop in creep resistance is observed.

It is also necessary that the γ/γ′ lattice misfit should conform to a small value, either positive or negative, since coherency is otherwise lost; thus limits are placed on its magnitude. The lattice misfit δ is defined as the mismatch between γ and γ′ phases, and is determined according to

$\begin{matrix} {\delta = \frac{2\left( {a_{\gamma^{\prime}} - a_{\gamma}} \right)}{a_{\gamma^{\prime}} + a_{\gamma}}} & (1) \end{matrix}$

where α_(γ)and α_(γ′) are the lattice parameters of the γ and γ′ phases.

Thus the model isolates all compositions in the design space which are calculated to result in a desired volume fraction of γ′, which have a lattice misfit γ′ of less than a predetermined magnitude.

In the fourth stage, merit indices are estimated for the remaining isolated alloy compositions in the dataset. These include: creep-merit index (which describes an alloy's creep resistance based solely on mean composition), strength-merit index (which describes an alloy's precipitation yield strength based solely on mean composition), density, cost, stable microstructure and gamma-prime solvus temperature.

In the fifth stage, the calculated merit indices are compared with limits for required behaviour, these design constraints are considered to be the boundary conditions to the problem. All compositions which do not fulfil the boundary conditions are excluded. At this stage, the trial dataset will be reduced in size quite markedly.

The final, sixth stage involves analysing the dataset of remaining compositions. This can be done in various ways. One can sort through the database for alloys which exhibit maximal values of the merit indices—the lightest, the most creep resistant, the most oxidation resistant, and the cheapest for example. Or alternatively, one can use the database to determine the relative trade-offs in performance which arise from different combination of properties.

The seven merit indicies are now described.

The first merit index is the creep-merit index. The overarching observation is that time-dependent deformation (i.e. creep) of a nickel-based superalloy occurs by dislocation creep with the initial activity being restricted to the γ phase. Thus, because the fraction of the γ′ phase is large, dislocation segments rapidly become pinned at the γ/γ′ interfaces. The rate-controlling step is then the escape of trapped configurations of dislocations from γ/γ′ interfaces, and it is the dependence of this on local chemistry in this case composition of the γ phase—which gives rise to a significant influence of alloy composition on creep properties.

A physically-based microstructure model can be invoked for the rate of accumulation of creep {dot over (ε)} strain when loading is uniaxial and along the

001

crystallographic direction. The equation set is

$\begin{matrix} {{\overset{.}{ɛ}}_{\langle 001\rangle} = {\frac{16}{\sqrt{6}}\rho_{m}\phi_{p}{D_{eff}\left( {1 - \phi_{p}} \right)}\left( {{1/\phi_{p}^{1/3}} - 1} \right)\sinh\left\{ \frac{{\sigma b}^{2}\omega}{\sqrt{6}K_{CF}{kT}} \right\}}} & (2) \\ {{\overset{.}{\rho}}_{m} = {C\overset{.}{ɛ}_{\langle 001\rangle}}} & (3) \end{matrix}$

where ρ_(m) is the mobile dislocation density, ϕ_(p) is the volume fraction of the γ′ phase, and ω is width of the matrix channels. The terms σ and T are the applied stress and temperature, respectively. The terms b and k are the Burgers vector and Boltzmann constant, respectively. The term K_(CF)=1+2ϕ_(p) ^(1/3)/3√{square root over (3π)}(1−ϕ_(p) ^(1/3)) is a constraint factor, which accounts for the close proximity of the cuboidal particles in these alloys. Equation 3 describes the dislocation multiplication process which needs an estimate of the multiplication parameter C and the initial dislocation density. The term D_(eff) is the effective diffusivity controlling the climb processes at the particle/matrix interfaces.

Note that in the above, the composition dependence arises from the two terms ϕ_(p) and D_(eff). Thus, provided that the microstructural architecture is assumed constant (microstructural architecture is mostly controlled by heat treatment) so that ϕ_(p) is fixed, any dependence upon chemical composition arises through D_(eff). For the purposes of the alloy design modelling described here, it turns out to be unnecessary to implement a full integration of Equations 2 and 3 for each prototype alloy composition. Instead, a first order merit index M_(creep) is employed which needs to be maximised, which is given by

$\begin{matrix} {M_{creep} = {\underset{i}{\Sigma}{x_{i}/{\overset{\sim}{D}}_{i}}}} & (4) \end{matrix}$

where x_(i) is the atomic fraction of solute i in the γ phase and {tilde over (D)}_(i) is the appropriate interdiffusion coefficient.

The second merit index is a strength merit index. For high nickel-based superalloys, the vast majority of strength comes from the precipitate phase. Therefore, optimising alloy composition for maximal precipitate strengthening is a critical design consideration. From hardening theory a merit index for strength, M_(strength), is proposed. The index considers the maximum possible precipitate strength determined to be the point where the transition from weakly coupled to strongly coupled dislocation shearing occurs which can be approximated using,

M _(strength) =M·1/2γ_(APB)Ø_(p) ^(1/2) /b   (5)

Where M is the Taylor factor, γ_(APB) is the anti-phase boundary (APB) energy, ϕ_(p) is the volume fraction of the γ′ phase and b is the Burgers vector.

From Equation 5 it is apparent that fault energies in the γ′ phase—for example, the anti-phase boundary APB energy—have a significant influence on the deformation behaviour of nickel-based superalloys. Increasing the APB energy has been found to improve mechanical properties including, tensile strength and resistance to creep deformation. The APB energy was studied for a number of Ni—Al—X systems using density functional theory. From this work the effect of ternary elements on the APB energy of the γ′ phase was calculated, linear superposition of the effect for each ternary addition was assumed when considering complex multicomponent systems, resulting in the following equation,

γ_(APB)=195−1.7x _(Cr)−1.7x _(Mo)+4.6x _(W)+27.1x _(Ta)+21.4x _(Nb)+15x _(Ti)   (6)

where, x_(Cr), x_(Mo), x_(W), x_(Ta), x_(Nb) and x_(Ti) represent the concentrations, in atomic percent, of chromium, molybdenum, tungsten, tantalum, niobium and titanium in the γ′ phase, respectively. The composition of the γ′ phase is determined from phase equilibrium calculations.

The third merit index is density. The density, ρ, was calculated using a simple rule of mixtures and a correctional factor, where, ρ_(i) is the density for a given element and x_(i) is the atomic fraction of the alloy element.

ρ=1.05[Σ_(i) x _(i)ρ_(i)]  (7)

The fourth merit index is cost. In order to estimate the cost of each alloy a simple rule of mixtures was applied, where the weight fraction of the alloy element, x_(i), was multiplied by the current (2017) raw material cost for the alloying element, c_(i).

Cost=Σ_(i)x_(i)c_(i)   (8)

The estimates assume that processing costs are identical for all alloys, i.e. that the product yield is not affected by composition.

A fifth merit index is based upon rejection of candidate alloys on the basis of unsuitable microstructural architecture made on the basis of susceptibility to TCP phases. To do this use is made of the d-orbital energy levels of the alloying elements (referred as Md) to determine the total effective Md level according to

M_(d) =Σ_(i)x_(i)Md_(i)   (9)

where the x_(i) represents the mole fraction of the element i in the alloy. Higher values of Md are indicative of higher probability of TCP formation.

A sixth merit index is the gamma-prime solvus temperature. The gamma-prime solvus is defined as the temperature where the volume fraction of gamma-prime tends to zero. This is determined using thermodynamic calculations—as previously described above in the second step of the Alloys-by-Design method. The phase diagram and thermodynamic properties for a specific alloy composition is calculated and used to find the temperature at which this phase transition occurs.

A seventh merit index is solid solution merit index. Solid solution hardening occurs in the (FCC) matrix phase which is referred to as gamma (γ), in particular this hardening mechanism is important at high temperatures for high strength and creep resistance. A model which assumes superposition of individual solute atoms on the strengthening of the matrix phase is employed. The solid solution strengthening coefficients, k_(i), for the elements considered in the design space: aluminium, cobalt, chromium, molybdenum, niobium, tantalum, titanium and tungsten are 225, 39.4, 337, 1015, 1183, 1191, 775 and 977 MPa/at. %^(1/2), respectively (H. Roth, C. Davis, and R. Thomson: Metallurgical and Materials Transactions A, 1997, vol. 28, pp. 1329-1335). The solid-solution index is calculated based upon the equilibrium composition of the matrix phase using the following equation,

M _(solid-solution)Σ_(i)(k _(i) ²√{square root over (x _(i))})   (10)

where, M solid-solution is the solid solution merit index and x_(i) is the concentration of element i in the γ matrix phase.

The ABD method described above was used to isolate the inventive alloy composition. The design intent for this alloy was to deliver substantially high strength, particularly high temperature strength and creep resistance in combination with a reasonable alloy cost, good oxidation/corrosion resistance and microstructural stability. The balance of properties for the new alloy enable higher maximum operating temperature in comparison to the prior art, particularly for use as a turbine disc application where the operating temperature for the component is 800° C. or greater.

The material properties—determined using the ABD method—for the nominal compositions of the commonly applied/researched alloys used for powder metallurgy (PM) turbine disc applications, listed in Table 1, are listed in Table 3. The design of the new alloy was considered in relation to the predicted properties listed for these alloys.

The rationale for the design of the new alloy is now described.

TABLE 3 Calculated phase fractions, misfit and merit indices made with the “Alloys-by-Design” software. Results for nickel-based superalloys listed in Table 1. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) N18 0.55 −0.61% 2.81 1472 8.1 17.7 0.94 1194 111 René88DT 0.34 0.04% 2.45 1392 8.5 16.7 0.91 1108 95 RR1000 0.41 −0.07% 2.77 1474 8.3 21.3 0.92 1139 94 ME3 0.48 −0.05% 3.2 1608 8.3 23.1 0.92 1163 93 Alloy10 0.52 0.01% 3.16 1616 8.4 21.0 0.91 1182 91 LSHR 0.51 −0.11% 3.5 1647 8.4 22.4 0.93 1171 96 U.S. Pat. No. 8,147,749 0.5 0.35% 3.32 1854 8.8 28.6 0.91 1141 89 N19 0.41 0.01% 2.41 1541 8.4 16.4 0.91 1136 100 U.S. Pat. No. 8,613,810 0.52 0.14% 3.22 1818 8.6 25.2 0.92 1156 94 US2015/0192002 0.51 0.17% 2.9 1719 8.6 21.6 0.92 1152 88

The additions of aluminium, niobium, tantalum and titanium influence the fraction of the precipitate hardening γ′ phase. A requirement for a strength merit index to be greater than 1900 MPa, representing a strength improvement relative to U.S. Pat. No. 8,147,749 the strongest alloy from the prior art listed, was chosen. It was found by plotting strength merit index vs volume fraction for all alloy compositions in the design space that with a gamma-prime volume fraction of greater than or equal to 51% at a temperature of 850° C. the desired strength could be achieved (FIG. 2).

FIG. 3 describes the relationship between alloy γ′ fraction (as contours) and alloy composition based upon aluminium plotted versus the sum of elements titanium, tantalum and niobium (according to the relationship 0.6W_(Ti)+0.31W_(Nb)+0.15W_(Ta)). The additions of titanium, niobium and tantalum have been given a weighting factor to convert the weight percent addition to an “aluminium equivalent”. Tantalum is included in this analysis as allowable levels of tantalum are quite high due to allowing the cost of the alloy to be quite high. This allows for direct comparison of the influence of elements which have very different densities. For example titanium has a density of 4.5 g/cm³ compared to aluminium with density of 2.7 g/cm^(3,) thus a factor of 0.6 is applied (i.e. 2.7/4.5=0.6). Similar to titanium, a constant is added to convert the elemental additions of niobium (8.57 g/cm³) and tantalum (16.4 g/cm³) to an “aluminium equivalent”, thus, niobium and tantalum have correctional factors (determined from their density relative to aluminium) of 0.31 and 0.15 respectively. From FIG. 3 it can be determined that to achieve an alloy with a γ′ fraction greater than or equal to 51% the additions of aluminium, titanium and niobium adhere to the following equation

f(γ′)=0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta)+0.94W _(Al)≥6.2   (11)

where f(γ′) is a numerical value and W_(Ti), W_(Nb), W_(Ta), and W_(Al) are the weight percent of titanium, niobium, tantalum and aluminium in the alloy respectively. To maintain easy ability to process the alloy, gamma-prime solvus is desirably less than 1180° C., this means it is beneficial to have a gamma-prime fraction of less than 62% (described later with reference to FIG. 4). To achieve a gamma-prime fraction of 62% or less the numerical value for f(γ′) must be 7.4 or less. Thus the numerical range for f(γ′) is preferably in the range between 6.2 and 7.4. It is more preferable for the alloy gamma-prime solvus should be less than 1170° C. This limits the gamma-prime fraction to 56% or less. To achieve that the numerical value for f(γ′) must be 6.7 or less.

The gamma-prime volume fraction is measured experimentally by the following procedure. After a substantially long thermal exposure at 850° C. the specimen is quenched in water and a section is taken through the material and polished using conventional/standard metallurgical preparation techniques for scanning electron microscopy. Once prepared the gamma/gamma-prime microstructure should be observable in a scanning electron microscope, particles of diameter 30 nm or greater should be observable. A minimum of 10 images are taken which provide a statistically representative dataset, the images should cover an area of at least 1 mm². The 2-dimensional images which reveal the gamma/gamma-prime microstructure should be processed to identify the gamma-prime phase, the area fraction of the gamma-prime phase should be measured. The area fraction of the phase is taken to be the volume fraction of gamma-prime and should lie between 51-62%.

In order to achieve a good balance between strength and ability to process the alloy it is beneficial if the solvus temperature of the γ′ phase is less than 1180° C. A γ′ solvus of less than 1180° C. is preferred as this allows for heat treatment above the γ′ solvus whilst reducing the susceptibility of the alloy to cracking on cool down from above the γ′ solvus temperature. Heat-treatment above the γ′ solvus temperature is desirable as this enables the growth of coarse grains which improves resistance to dwell fatigue, this damage mechanism is often a life limiting factor in this class of alloys. From FIG. 4 a relationship between alloy γ′ solvus and alloy composition, based upon aluminium and the sum of elements titanium, tantalum and niobium (according to the relationship 0.6W_(Ti)+0.31W_(Nb)+0.15W_(Ta)) was determined. Additions of aluminium, titanium and niobium adhere to the following equation

f(solvus)=1.1W _(Al)+0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta)   (12)

where f(solvus) is a numerical value. In order to produce an alloy with a solvus of 1180° C. or less the value for f(solvus) should be less than 8.0. This restriction limits gamma-prime volume fraction to 62%, see FIG. 3. Preferably the numerical value for f (solvus) should be less than 7.3 to produce an alloy with a solvus less than 1170° C. which will further improve the ability to process the alloy. Therefore it is preferable to limit the gamma-prime fraction to 56%, see FIG. 3.

It is beneficial to select alloys where the ratio of the elements satisfies the following relationship 1.0≤(0.6W_(Ti)+0.31W_(Nb)+0.15W_(Ta))/W_(Al)≤1.3 (see D. J. Crudden, N. Warnken, A. Mottura, and R. C. Reed. Modelling of the influence of alloy composition on flow stress in high-strength nickel-based superalloys. Acta Materialia, 7: 356-370, 2014). The minimum limit of 1.0 is desirable as this enable high alloy strengths to be achieved due to substantial increases in APB energy, see Equation (6). High APB energies are desirable as this promotes resistance to precipitate shearing imparting strength in the alloy in terms of both tensile strength and creep strength, particularly in intermediate temperature regimes between 600-850° C. which is the intended operating temperature range for this alloy. The maximum limit of 1.3 is preferred to maintain stability of the γ′phase. At values beyond this the formation of unwanted eta phase (Ni₃Ti) and delta phase (Ni₃Nb) are likely to occur, these phases reduce alloy ductility and fatigue resistance. Niobium must also be limited to less than 4.0 wt. % as high concentrations of niobium stabilise the niobium rich delta phase Ni₃Nb. Alloys where the ratio (0.6W_(Ti)+0.31 W_(Nb)+0.15 Wm)/W_(Al) lies between 1.0-1.3 will have a desirable combination of strength and gamma-prime stability. It is more preferable that the alloy contains less than 3 wt. % niobium as niobium has been found to increase rates of dwell crack propagation in these alloys due to the formation of niobium oxides. Most preferably the niobium content is limited to less than 2 wt. % as this inhibits the formation of niobium oxides and enables better resistance to dwell crack propagation.

FIG. 3 describes the relationship between alloy γ′ fraction and aluminium and the sum of elements titanium, tantalum and niobium (according to the relationship 0.6W_(Ti)+0.31W_(Nb)+0.15W_(Ta)). Delineated on this figure are the relationships for gamma-prime solvus where f(solvus) is 8.0 and 7.3. Also delineated is the relationships where (0.6W_(Ti)+0.31W_(Nb)+0.31W_(Ta))/W_(Al) is 1.0 and 1.3. The hatched area on FIG. 3 denotes the preferred compositional space in this invention. From the Figure it can be seen that in order to have a gamma-prime fraction of greater than 51% an aluminium content of greater than 2.8 wt. % is necessary (FIG. 3). To limit the gamma-prime solvus to less than 1180° C. it is necessary to limit the aluminium content to 3.7 wt. % or less. Preferably the aluminium content is limited to 3.5 wt. % as this helps achieve a lower gamma-prime solvus temperature. More preferably aluminium is limited to 3.1 wt. % or less as this achieves an even better combination of strength and lowering of gamma-prime solvus temperatures.

FIGS. 5-9 show the influence of titanium, niobium and tantalum on alloy strength, in terms of strength merit index when the aluminium concentration is restricted between 2.8 and 3.7 wt. %. On each figure the limit for gamma-prime solvus (f(solvus)=8.0) is plotted when the maximum concentration of aluminium (3.7 wt. %) is included in the alloy. From FIGS. 5-9 it is possible to determine a relationship between alloy chemistry and predicted strength. Additions of the elements adhere to the following equation when the aluminium content is between 2.8 and 3.7 wt. %

f(strength)=0.6W _(Ti)+0.44W _(Nb)+0.27W _(Ta)   (13)

where f(strength) is a numerical value which must be 4.2 or greater in order to achieve an alloy with a strength merit index of greater than 1900 MPa. The constants for niobium and tantalum have been derived based on Equation 5 by normalising their strengthening effect to titanium and also through consideration of their “aluminium equivalence”. For example niobium has an aluminium equivalence of 0.31 and it has an effect on APB energy which is 1.4 times greater than titanium, thus (0.31×1.4) a factor of 0.44 is applied. For tantalum the aluminium equivalence is 0.15 and the effect on APB energy is 1.8 times greater than titanium, thus (0.15×1.8) a factor of 0.27 is applied.

From FIGS. 5-9 it is clear that niobium and titanium alone cannot easily achieve the required level of strength merit index (1900 MPa) whilst meeting the preferred low solvus temperature. Therefore the alloy of this invention must include a tantalum content of greater than 2.4 wt. %. Higher levels of tantalum enable higher strength with lower γ′ solvus then is possible with just titanium and niobium levels and therefore it is preferable that the tantalum content of the alloy is greater than 2.45 wt. % or 2.5 wt. %. It is most preferable that the tantalum is greater than or equal to 3.4 wt. %. Even more preferably, particularly when niobium is limited to 2.0 wt. %, tantalum content in the alloy is greater than 5.1 wt. % as this achieves a higher alloy strength, resulting in a strength merit index of 1950 MPa or greater. Preferably niobium is present in an amount of at least 1.0 wt %, to contribute to the strength of the alloy.

The titanium content of the alloy is limited to less than 5.1 wt. % to ensure that the alloy has the correct balance of strength (in terms of strength merit index greater than 1900 MPa) and ability to be manufactured (in terms of gamma-prime solvus of less than 1180° C.), see FIG. 5. Additions of titanium to the alloy have the effect of increasing alloy strength. However, titanium may have a negative influence on oxidation, therefore it is preferable that titanium is limited to 4.8 wt. % (FIG. 6), more preferably 4.6 wt. % (FIG. 7) so that a balance an improved balance of strength and oxidation resistance is achieved. Even more preferably titanium content can be limited to 4.2 wt. %, particularly when the tantalum content is 5.1 wt. %, whilst still achieving a higher strength of 1950 MP, this provides an even better combination of alloy strength and oxidation resistance

Based on maximum tantalum concentration of 9.5 wt. % (described later in terms of alloy density using FIG. 11) and maximum concentration of niobium (4 wt. %), the titanium content of the alloy is at least 0.3 wt. % based upon the relationship described for f(strength). It is preferable to have titanium content of at least 1.1 wt. %, particularly when niobium is limited to 3.0wt. % (to improve stability). It is more preferable to have a titanium content of at least 1.8 wt. %, particularly when niobium content of less than 2.0 wt. % (further to improve stability) is used. It is desirable also to limit titanium content in the alloy due to its negative impact on oxidation.

Along with increased strengthening contribution from precipitate hardening resulting from the gamma-prime phase (calculated in terms of strength merit index) it is desirable to design an alloy which has a strong gamma matrix phase. The strength of the gamma matrix can be calculated in terms of a solid solution merit index. The solid solution strengthening is particularly important for imparting high temperature strength (both tensile and creep strength) in the alloy. The solid solution strengthening of the gamma phase of the alloy is mainly dependent upon the additions of elements tungsten and molybdenum, as the coefficients for these elements is large and they strongly partition to the gamma-phase, unlike niobium and tantalum which also have large strengthening coefficients but they have limited partitioning to the gamma phase-see Equation 10. FIG. 10 shows the relationship between molybdenum, tungsten and solid solution index. A coefficient of 0.5 is applied to the tungsten content as it is approximately twice the density of molybdenum; this factor accounts for differences in density. To provide improved creep strength over the prior art—in particular alloys in Table 3 which have a strength merit index of greater than 1800 MPa it is desirable for the solid solution index to be greater than 100 MPa. From FIG. 10 it is possible to determine a relationship between additions of tungsten and molybdenum and solid solution index

f(solid solution)=1.29W _(Mo)+0.5W _(W)

where f(solid solution) is a numerical value which must be 5.7 or greater and W_(Mo), and W_(W) are the weight percent of molybdenum and tungsten in the alloy respectively in order to achieve an alloy with a solid solution merit index of greater than 100 MPa. Due to the density constraint of 8.6 g/cm³ which is imposed, tungsten is limited to 8.0 wt. %, preferably 7.8% and most preferably 6.9 wt. %, described in following section with reference to FIG. 11. Therefore based on the upper limit for tungsten (8.0 wt. %) the alloy contains a minimum of 1.1 wt. % molybdenum to satisfy the solid solution requirements. Molybdenum is preferably present greater than 1.2 wt. %, particularly when tungsten is limited to 8.2 wt. %. More preferably the molybdenum is more preferably greater than 1.7 wt. % particularly when tungsten is limited to 6.9 wt. %. Desirably f(solid solution) is 6.0 or greater, achieving a solid solution merit index of greater than 103 MPa.

In combination with an increase in alloy strength it is beneficial to isolate alloys which have a good combination of strength and weight. Therefore it is desirable to control the density; the design target for this alloy was to limit the density to 8.6 g/cm³ or less, resulting in a density equivalent to high strength alloys in Table 3. FIG. 11 shows the effect of tantalum and tungsten on alloy density, these alloying elements have the strongest influence on alloy density as they are substantially heavier than nickel (8.9 g/cm³); tungsten and tantalum have densities of 19.3 g/cm³ and 16.4 g/cm³ respectively. From FIG. 11 it is seen that the tantalum content of the present invention should be limited to 9.5 wt. % or less, to limit the density of the alloy. From FIG. 11 it is possible to determine a relationship between additions of tungsten and tantalum and alloy density, and additions of the elements adhere to the following equation

f(density)=W _(Ta)+0.88W _(W)

where f(density) is a numerical value which must be less than 9.4 in order to achieve density of less than 8.6 g/cm³. Based upon the minimum required level of tantalum (>2.4 wt. %) it is desirable to limit the tungsten content to 8.0 wt. %. Preferably tungsten is limited to 7.8 wt. % or 6.9 wt. % to keep the density of the alloy low and allowing greater levels of tantalum for increased strength. It is also preferable to have a tungsten content of greater than 3.3 wt. %, described later in terms of creep merit index and alloy cost, therefore it is preferable to limit the tantalum content to 6.4 wt. % or less to ensure a good balance of creep resistance, alloy cost and density.

The additions of cobalt and tantalum most strongly effect the alloy cost. FIG. 12 shows the relationship between tantalum and cobalt additions and alloy cost. From FIG. 1 it is determined that it is desirable to limit alloy cost to 26 $/kg, as any further increases in strength require a substantial increase in cost. Moreover this results in a cost reduction relative to other high strength alloys listed in Table 3. Therefore, a cost limit of 26 $/kg provides an attractive balance of alloy cost and performance. The cobalt concentration of the alloy is limited to 24.9 wt. %, based upon the minimum tantalum content of the alloy (>2.4 wt. %). The relationship between alloy chemistry and predicted cost is as follows

f(cost)=W _(Ta)+0.47W _(Co)

where f(cost) is a numerical value which must be less than 14.1 to achieve cost of 26 $/kg or less. With a minimum tantalum content of 2.4 wt. %, this gives a maximum cobalt concentration of 24.9 wt. %. Prefereably tantalum content is greater than 2.45 wt. % or 2.5 wt. % to provide a higher level of strength. It is prefered to limit cobalt to 24.6 wt. %. More preferably tantalum should be greater than 3.4 wt. %. It is more preferable that cobalt is limited to 22.8 wt. %. Most preferably tantalum is greater than 5.1 wt. %. Most preferably cobalt is limited to 19.1 wt. %. Even more preferably a cost limit of 22 $/kg, therefore f(cost) should be less than 10.3, therefore it is preferable to limit the cobalt concentration of the alloy to 16.8 wt. %. Particularly if tantalum concentration is greater than 3.4 wt. %, it is more preferable that cobalt is limited to 14.7 wt. %. Particularly if tantalum concentration is greater than 5.1 wt. % most preferably cobalt is limited to 11.1 wt. %

Creep resistance in terms of creep merit index is dependent upon additions of molybdenum, tungsten and cobalt (FIGS. 13-17). A target creep merit index of 2.8×10⁻¹⁶ m⁻²s is desired to provide sufficient creep resistance to the alloy. From FIGS. 13-17 it is possible to determine a relationship between additions of tungsten, cobalt and molybdenum on creep merit index. Additions of the elements adhere to the following equation

f(creep)=W _(W)+0.93W _(Co)+0.84W _(Mo)

where f(creep) is a numerical value which must be greater than 17.2 in order to achieve an alloy with a creep merit index of greater than 2.8×10⁻¹⁶ m⁻²s. Based upon the upper limit for molybdenum (4.4 wt. %) and tungsten (8.0 wt. %) it is necessary to have a least 5.9 wt. % cobalt in the alloy to achieve a creep merit index of at least 2.8×10⁻¹⁶ m⁻²s. Preferably tungsten is limited to 7.8 wt. % and in this case it is preferable that cobalt is at least 6.1 wt. % which is desirable in any case as it increases creep strength. More preferably tungsten is limited to 6.9 wt. % and in this case it is more preferable to have a cobalt concentration of 7.1 wt. %. As described in relation to alloy cost it is most preferable to reduce the use of cobalt in the alloy, preferably to less than 11.1 wt. %. It is preferable to have a tungsten content of at least 3.2 wt. %.

FIGS. 18-23 describe the effect of chromium, tungsten and molybdenum on the stability number. A higher stability number results in an alloy which is more prone to TCP phase formation. Limiting or stopping the precipitation of TCP phase formation is beneficial as these phases lead to deterioration in material properties over time. A chromium level of greater than 9.0 wt. % is desirable in order to achieve a good level of oxidation resistance as this level of chromium will allow formation of a protective chromia oxide scale. A stability target of less than 0.92 in order to ensure microstructural stability and avoid TCP formation, see prior art alloys in Table 3. More preferably a stability target of less than 0.91 is desirable in order to ensure better microstructural stability and avoid TCP formation. From FIGS. 18-23 it is determined that the additions of the elements molybdenum tungsten and chromium adhere to the following equation

f(stability)=W _(W)+1.62W _(Cr)+2.0W _(Mo)

where f(stability) is a numerical value which must be less than 30.2 in order to achieve an alloy with a stability number of less than 0.92. Through consideration of the equation for solid solution strengthening where f(solid solution)>5.7 it is determined that if a maximum limit to molybdenum content is set to 4.4 wt. %, no tungsten may be present. Thus the maximum molybdenum content is 4.4 wt. %. By limiting molybdenum content to 4.4 wt. % it possible to get an excellent balance of high temperature strength (gained from solid solution strengthening) and oxidation resistance as chromium content can be maximised whilst satisfying the requirement for stability number. Thus based on the maximal molybdenum content of 4.4 wt. % the chromium content is limited to 13.2 wt. % to ensure a good balance of high temperature strength and oxidation resistance.

It is preferable to limit the stability number to 0.91. To do this the numerical value for f(stability) should be less 28.3. Therefore it is preferable to limit chromium content to 11.9 wt. % as this will limit the stability number to 0.91 providing better microstructural stability based upon previously described molybdenum and tungsten content. It is more preferable to limit the chromium content to 11.5 wt. % for the best balance of high temperature strength, cost and stability, as the stability number can be less than 0.92. Even more preferably it is more preferable to limit chromium content to 11.2 wt. % as this can limit the Md number 0.91.

Additions of carbon, boron and zirconium are required in order to provide strength to grain boundaries. This is particularly beneficial for the creep and fatigue properties of the alloy. Carbon is added to act as a grain boundary pinning particle, this is necessary when heat treatment is conducted above the gamma-prime solvus temperature to inhibit excessive grain growth. The carbon concentrations should range between 0.01 wt. % and 0.1 wt. %. Preferably, the levels of carbon are between 0.2 and 0.06, this range provides a better distribution of carbide phases for controlling alloy microstructure, in particular grain size.

The boron concentration should range between 0.001 and 0.1 wt. %. The addition of boron can improve creep ductility and grain boundary strength through the formation of boride phases. Preferably the boron content in the alloy is between 0.01 and 0.05 wt. %, as this provides a desirable level of the boride phase.

The zirconium concentrations should range between 0.001 wt. % and 0.3 wt. %, preferably between 0.02 and 0.1 wt. %. Zirconium plays a role in guttering unwanted impurities in the ally, for example, oxygen and sulphur. These impurities may lead to embrittlement of the alloy particularly due to grain boundary embrittlement.

It is beneficial that when the alloy is produced, it is substantially free from incidental impurities. These impurities may include the elements sulphur (S), manganese (Mn) and copper (Cu). The element sulphur should remain below 0.003 wt. % (30 PPM in terms of mass). Manganese is an incidental impurity which is limited to 0.25 wt. %, preferably this limited to less than 0.1 wt. %. Copper (Cu) is an incidental impurity which is preferably limited to 0.5 wt. %. The presence of Sulphur above 0.003 wt. %, can lead to embrittlement of the alloy and sulphur also segregates to alloy/oxide interfaces formed during oxidation, preferably sulphur levels of less than less than 0.001 wt. %. Vanadium is an incidental impurity, vanadium negatively influences the oxidation behaviour of the alloy and is which is preferably limited to 0.5 wt. %, preferably less than 0.3 wt. % and most preferably this limited to less than 0.1 wt. %. This segregation may lead to increased spallation of protective oxide scales. If the concentrations of these incidental impurities exceed the specified levels, issues surrounding product yield and deterioration of the material properties of the alloy is expected.

Additions of hafnium (Hf) of up to 0.5 wt. %, are beneficial for tying up incidental impurities in the alloy and also for providing strength. Hafnium is a strong carbide former it can provide additional grain boundary strengthening. More preferably hafnium is limited to 0.2 wt. %, more preferably less than 0.1 wt. % as the elemental cost is significant, additions have a negative impact on alloy cost.

Additions of the so called ‘reactive-elements’, Yttrium(Y), Lanthanum (La) and Cerium (Ce) may be beneficial up to levels of 0.1 wt. % to improve the adhesion of protective oxide layers, such as Cr₂O₃. These reactive elements can ‘mop-up’ tramp elements, for example sulphur, which segregates to the alloy oxide interface weakening the bond between oxide and substrate leading to oxide spallation. Additions of Silicon (Si) up to 0.5 wt. % may be beneficial, it has been shown that additions of silicon to nickel based superalloys at levels up to 0.5 wt. % are beneficial for oxidation properties. In particular silicon segregates to the alloy/oxide interface and improves cohesion of the oxide to the substrate. This reduces spallation of the oxide, hence, improving oxidation resistance.

In an embodiment of the invention is desirable to include iron in substitution for nickel content. This has the benefit of reducing alloy cost and increasing the ability for the alloy to be recycled. Additions of iron may result in increased microstructural instability. Limiting iron additions to a level of 4.0 wt. % produces a good balance of low cost, improved recyclability and microstructural stability, more preferably a range between 1.0 wt. % and 2.0 wt. % is desirable.

Based upon the description of the invention presented in this section the broad range for the invention is listed in Table 4. A preferable range is also given in Table 4.

TABLE 4 Compositional range in wt. % for the newly design alloy. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb Fe Min 9.0 5.9 1.1 0.0 2.8 0.3 >2.4 0.0 0.0 Max 13.2 24.9 4.4 8.0 3.7 5.1 9.5 4.0 4.0 Preferable Min 9.0 6.1 1.2 3.2 2.8 1.1 2.5 0.0 0.0 Preferable Max 11.9 14.7 4.4 7.8 3.7 4.6 9.5 3.0 2.0 Most Preferable 9.0 7.9 1.7 3.2 2.8 1.8 5.1 1.0 1.0 Min Most Preferable 11.5 11.1 4.4 6.9 3.5 4.2 6.4 2.0 2.0 Max

The following Section describes example compositions for the present invention. The calculated properties for these new alloys are listed. The rationale for the design of these alloys is now described.

EXAMPLES OF THE INVENTION

Two examples of the invention are described in Table 5. The predicted properties of the alloys are listed in Table 6. The alloys listed have the benefit of increased strength in terms of strength merit index—over the prior art alloys listed in Table 3. This improvement in strength is achieved in combination with a solvus temperature which is equivalent to the prior art to enable manufacture of the alloy. The density of the alloys in Table 5 is also equivalent to alloys listed in Table 3 which have a strength index of greater than 1700 MPa. These improvements in material performance are attained whilst maintaining equivalent or lower cost than the alloys which have a strength index of greater than 1700 MPa.

TABLE 5 Examples of alloys in the present invention Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD3 11.50 8.00 2.00 6.80 3.00 4.00 3.00 3.00 0.050 0.025 0.060 0.00 ABD-PMD4 10.00 10.00 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00

TABLE 6 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software. Results for compositions listed in Table 5. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD3 0.58 −0.01% 2.8 1970 8.6 18.5 0.92 1163 105 ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107

In Table 7 the quantity of gamma-prime forming elements is modified in alloy ABD-PMD4. Alloys ABD-PMD5-ABD-PMD8 have a lower level of gamma-prime forming elements, this results in a reduction of alloy strength, through lowering of the gamma-prime volume fraction (Table 8). However, a higher strength than prior art in Table 3 is still achieved. Reducing the gamma-prime forming elements has the benefit of reducing gamma-prime solvus, this improves the ability to form the alloy through thermo-mechanical processing. A lower solvus also makes the alloy less sensitive to cracking during heat-treatment. Example alloy ABD-PMD10 has a higher level of gamma-prime forming elements in comparison to ABD-PMD4. It is seen that strength can be increased further whilst still maintaining a gamma-prime solvus of less than 1180° C. This alloy has greater mechanical strength than ABD-PMD4 while remaining amenable to thermo-mechanical but does not quite achieve the preferred gamma-prime solvus of less than 1170° C.

TABLE 7 Example compositions where the gamma-prime forming element content has been modified in example alloy ABD-PMD4 to modify gamma-prime content and gamma-prime solvus. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD4 10.00 10.50 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD5 10.00 10.50 3.00 4.90 2.82 3.76 4.70 1.88 0.050 0.030 0.060 0.00 ABD-PMD6 10.00 10.50 3.00 4.90 2.88 3.83 4.79 1.92 0.050 0.030 0.060 0.00 ABD-PMD7 10.00 10.50 3.00 4.90 2.92 3.89 4.86 1.94 0.050 0.030 0.060 0.00 ABD-PMD8 10.00 10.50 3.00 4.90 2.96 3.95 4.93 1.97 0.050 0.030 0.060 0.00 ABD-PMD9 10.00 10.50 3.00 4.90 3.00 4.01 5.01 2.00 0.050 0.030 0.060 0.00 ABD-PMD10 10.00 10.50 3.00 4.90 3.05 4.07 5.08 2.03 0.050 0.030 0.060 0.00

TABLE 8 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software for alloys where the gamma-prime forming element content has been modified in example alloy ABD-PMD4. Results for compositions listed in Table 7. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107 ABD-PMD5 0.53 0.09% 2.8 1900 8.6 21.4 0.92 1154 104 ABD-PMD6 0.55 0.07% 2.8 1927 8.6 21.6 0.92 1159 105 ABD-PMD7 0.56 0.05% 2.8 1948 8.6 21.6 0.92 1162 106 ABD-PMD8 0.57 0.02% 2.8 1970 8.6 21.7 0.93 1165 106 ABD-PMD9 0.58 0.00% 2.9 1991 8.6 21.8 0.93 1168 107 ABD-PMD10 0.59 −0.03% 2.9 2013 8.6 21.9 0.93 1171 108

In Table 9 titanium and tantalum have been varied keeping the value for 0.6W_(Ti)+0.15W_(Ta) constant. Alloys ABD-PMD11-PMD13 have lower tantalum contents relative to PMD-4, titanium is used in substitution for tantalum. This is particularly advantageous when there is a need to reduce the elemental cost of the alloy. In Alloys PMD-14-PMD-16 tantalum has been substituted for titanium. This is advantageous as it further increases the alloy strength and decreases the gamma-prime solvus improving the ability to manufacture the alloy. The reduction in titanium is also beneficial for improving the oxidation resistance, these improvements come with an increase in alloy cost.

TABLE 9 Example compositions where the elements titanium and tantalum have been switched in equal proportions in example alloy ABD-PMD4. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD4 10.00 10.50 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD11 10.00 10.50 3.00 4.90 3.00 4.60 2.60 2.00 0.050 0.030 0.060 0.00 ABD-PMD12 10.00 10.50 3.00 4.90 3.00 4.40 3.40 2.00 0.050 0.030 0.060 0.00 ABD-PMD13 10.00 10.50 3.00 4.90 3.00 4.20 4.20 2.00 0.050 0.030 0.060 0.00 ABD-PMD14 10.00 11.20 3.39 3.90 3.00 3.80 5.80 2.00 0.050 0.030 0.060 0.00 ABD-PMD15 10.00 12.00 3.78 2.90 3.00 3.60 6.60 2.00 0.050 0.030 0.060 0.00 ABD-PMD16 10.00 12.70 4.16 1.90 3.00 3.40 7.40 2.00 0.050 0.030 0.060 0.00

TABLE 10 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software for alloys where elements titanium and tantalum have been switched in equal proportions in example alloy ABD-PMD4. Results for compositions listed in Table 9. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107 ABD-PMD11 0.57 0.00% 2.8 1900 8.5 19.0 0.92 1180 104 ABD-PMD12 0.57 0.00% 2.8 1928 8.5 19.9 0.92 1175 105 ABD-PMD13 0.57 0.00% 2.8 1958 8.5 20.9 0.92 1173 106 ABD-PMD14 0.57 0.02% 2.9 2013 8.6 23.1 0.92 1163 107 ABD-PMD15 0.57 0.04% 2.9 2038 8.6 24.4 0.92 1159 107 ABD-PMD16 0.57 0.06% 2.9 2062 8.6 25.7 0.92 1154 107

In Table 11 titanium and niobium have been varied keeping the value for 0.6 W_(Ti)+0.31 Nb constant. Alloys ABD-PMD17-PMD20 have lower niobium contents relative to PMD-4, titanium is used in substitution for niobium. This is particularly advantageous when there is a need to reduce the niobium content in the alloy as reducing niobium may result in a lower susceptibility to the formation of niobium oxides which can accelerate oxidation assisted cracking mechanism in nickel superalloys. In Alloys PMD-21-PMD-24 niobium has been substituted for titanium. This is advantageous as it further increases the alloy strength and decreases the gamma-prime solvus improving the ability to manufacture the alloy. The reduction in titanium is also beneficial for improving the oxidation resistance, these improvements come with an increase in alloy cost.

TABLE 11 Example compositions where the elements titanium and niobium have been switched in equal proportions in example alloy ABD-PMD4. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD4 10.00 10.50 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD17 10.00 10.50 3.00 4.90 3.00 4.80 5.00 0.45 0.050 0.030 0.060 0.00 ABD-PMD18 10.00 10.50 3.00 4.90 3.00 4.60 5.00 0.84 0.050 0.030 0.060 0.00 ABD-PMD19 10.00 10.50 3.00 4.90 3.00 4.40 5.00 1.23 0.050 0.030 0.060 0.00 ABD-PMD20 10.00 10.50 3.00 4.90 3.00 4.20 5.00 1.61 0.050 0.030 0.060 0.00 ABD-PMD21 10.00 10.50 3.00 4.90 3.00 3.80 5.00 2.39 0.050 0.030 0.060 0.00 ABD-PMD22 10.00 10.50 3.00 4.90 3.00 3.60 5.00 2.77 0.050 0.030 0.060 0.00 ABD-PMD23 10.00 10.50 3.00 4.90 3.00 3.40 5.00 3.16 0.050 0.030 0.060 0.00 ABD-PMD24 10.00 10.50 3.00 4.90 3.00 3.20 5.00 3.55 0.050 0.030 0.060 0.00

TABLE 12 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software for alloys where elements titanium and niobium have been switched in equal proportions in example alloy ABD-PMD4. Results for compositions listed in Table 11. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107 ABD-PMD17 0.58 −0.05% 2.8 1937 8.5 21.3 0.92 1177 106 ABD-PMD18 0.58 −0.03% 2.8 1950 8.5 21.5 0.92 1175 106 ABD-PMD19 0.58 −0.02% 2.8 1964 8.6 21.6 0.92 1173 106 ABD-PMD20 0.58 −0.01% 2.8 1977 8.6 21.7 0.92 1170 107 ABD-PMD21 0.58 0.01% 2.9 2002 8.6 21.9 0.92 1165 108 ABD-PMD22 0.58 0.02% 2.9 2014 8.6 22.0 0.92 1163 108 ABD-PMD23 0.58 0.03% 2.9 2025 8.6 22.2 0.92 1160 109 ABD-PMD24 0.57 0.03% 2.9 2036 8.6 22.3 0.92 1157 109

In Table 13 the tungsten level in the alloy ABD-PMD-4 has been increased to improve the creep merit index resulting in improved alloy creep resistance. The molybdenum level in the alloy has been controlled to maintain alloy stability and solution strengthening. Tantalum and titanium have been substituted to maintain a high alloy strength and a low alloy density. The reduction in tantalum also has the benefit of reducing alloy cost.

TABLE 13 Example compositions where the elements molybdenum and tungsten have been switched in equal proportions in example alloy ABD-PMD4. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD4 10.00 10.50 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD25 10.00 10.50 2.80 5.42 3.00 4.11 4.55 2.00 0.050 0.030 0.060 0.00 ABD-PMD26 10.00 10.50 2.60 5.93 3.00 4.23 4.09 2.00 0.050 0.030 0.060 0.00 ABD-PMD27 10.00 10.50 2.40 6.45 3.00 4.34 3.64 2.00 0.050 0.030 0.060 0.00 ABD-PMD28 10.00 10.50 2.20 6.96 3.00 4.45 3.18 2.00 0.050 0.030 0.060 0.00 ABD-PMD29 10.00 10.50 2.00 7.48 3.00 4.57 2.73 2.00 0.050 0.030 0.060 0.00

TABLE 14 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software for alloys where elements molybdenum and tungsten have been switched in equal proportions in example alloy ABD-PMD4. Results for compositions listed in Table 13. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107 ABD-PMD25 0.58 −0.01% 2.9 1977 8.6 21.3 0.92 1171 107 ABD-PMD26 0.58 −0.02% 2.9 1965 8.6 20.7 0.92 1173 107 ABD-PMD27 0.58 −0.02% 3.0 1953 8.6 20.2 0.92 1175 107 ABD-PMD28 0.58 −0.03% 3.0 1941 8.6 19.7 0.92 1177 106 ABD-PMD29 0.58 −0.04% 3.1 1930 8.6 19.1 0.92 1179 106

In Table 15 the levels of molybdenum and tungsten have been reduced and replaced with chromium in order to improve the alloys creep resistance. Cobalt levels in the alloy are increased to maintain a high creep merit index for good creep resistance.

TABLE 15 Example compositions where the elements molybdenum and tungsten have been lowered and chromium content has been increased in example alloy ABD-PMD4. Alloy (wt. %) Cr Co Mo W Al Ti Ta Nb C B Zr Hf ABD-PMD4 10.00 10.50 3.00 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD30 10.16 10.60 2.87 4.90 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD31 10.38 10.90 2.74 4.80 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD32 10.57 11.00 2.61 4.75 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD33 11.00 11.00 2.61 4.75 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD34 11.50 11.00 2.61 4.75 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00 ABD-PMD35 12.00 11.00 2.61 4.75 3.00 4.00 5.00 2.00 0.050 0.030 0.060 0.00

TABLE 16 Calculated phase fractions and merit indices made with the “Alloys-by-Design” software for alloys where the elements molybdenum and tungsten have been lowered and chromium content has been increased in example alloy ABD-PMD4. Results for compositions listed in Table 15. Alloy (wt. %) Strength Solid γ/γ′ Creep Merit Merit Solution Misfit Index Index Density Cost γ′ Solvus Index γ′ (%) (m⁻²s × 10⁻¹⁶) (Mpa) (g/cm³) ($/kg) Md_(γ) (° C.) (Mpa) ABD-PMD4 0.58 0.00% 2.9 1989 8.6 21.8 0.92 1168 107 ABD-PMD30 0.58 0.02% 2.9 1990 8.6 21.8 0.92 1168 106 ABD-PMD31 0.58 0.05% 2.9 1991 8.6 22.0 0.92 1167 105 ABD-PMD32 0.58 0.07% 2.9 1992 8.6 22.0 0.92 1167 103 ABD-PMD33 0.58 0.05% 2.9 1996 8.6 22.0 0.92 1165 104 ABD-PMD34 0.58 0.03% 2.9 1999 8.6 22.0 0.92 1162 104 ABD-PMD35 0.58 0.01% 2.9 2003 8.6 22.0 0.92 1159 104 

1. A nickel-based alloy composition consisting, in weight percent, of: 9.0 to 13.2% chromium, 5.9 to 24.9% cobalt, 0.0 to 4.0% iron, 1.1 to 4.4% molybdenum, 0.0 to 8.0% tungsten, 2.8 to 3.7% aluminium, 0.3 to 5.1% titanium, 0.0 to 4.0% niobium, more than 2.4% tantalum and 9.5% or less tantalum, 0.01 to 0.1% carbon, 0.001 to 0.1% boron, 0.001 to 0.3% zirconium, 0.0 to 0.5% silicon, 0.0 to 0.1% yttrium, 0.0 to 0.1% lanthanum, 0.0 to 0.1% cerium, 0.0 to 0.003% sulphur, 0.0 to 0.25% manganese, 0.0 to 0.5% vanadium, 0.0 to 0.5% copper, and 0.0 to 0.5% hafnium, the balance being nickel and incidental impurities, wherein the following equations are satisfied in which W_(Mo), W_(w), W_(Nb), W_(Ta) and W_(Ti) are the weight percent of molybdenum, tungsten, niobium, tantalum and titanium in the alloy respectively 29W _(Mo)0+0.5W _(w)≥5.7 0.6W _(Ti)+0.44W _(Nb)+0.27W _(Ta)≥4.2.
 2. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(Nb), W_(Ta), W_(Ti) and W_(Al) are the weight percent of niobium, tantalum, titanium and aluminium in the alloy respectively 6.2≤0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta)+0.94W _(Al)≤7.4 preferbly 6.2≤0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta)+0.94W _(Al)≤6.7
 3. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(Nb), W_(Ta), W_(Ti) and W_(Al) are the weight percent of niobium, tantalum, titanium and aluminium in the alloy respectively 1.0≤(0.6W _(Ti)+0.31W _(Nb)+0.15W _(Ta))/W _(Al)≤1.3
 4. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(W) and W_(Mo) are the weight percent of tungsten and molybdenum, in the alloy respectively 6.0≤1.29W _(Mo)+0.5W _(W)
 5. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(Ta) and W_(W) are the weight percent of tantalum and tungsten, in the alloy respectively W _(Ta)+0.88W _(W)≤9.4
 6. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(Ta) and W_(Co) are the weight percent of tantalum and cobalt, in the alloy respectively W _(Ta)+0.47W _(Co)≤14.1
 7. The nickel-based alloy composition of claim 1, wherein the following equation is satisfied in which W_(W), W_(Mo) and W_(Cr) are the weight percent of tungsten and chromium and molybdenum, in the alloy respectively W _(W)+1.62W _(Cr)+2.0W _(Mo)≤30.2 preferably W _(W)+1.62W _(Cr)+2.0W _(Mo)≤28.3
 8. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 3.5 wt. % or less aluminium, preferably 3.1 wt. % or less aluminium.
 9. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 2.45 wt. % or more tantalum, preferably 2.5 wt. % or more tantalum, more preferably 3.4 wt. % or more tantalum, even more preferably 5.1 wt. % or more tantalum.
 10. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 6.4 wt. % or less tantalum
 11. (canceled)
 12. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 3.0 wt. % or less niobium, preferably 2.0 wt. % or less niobium.
 13. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 1.1 wt. % or more of titanium, preferably 1.8 wt. % or more titanium.
 14. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 4.8 wt. % or less titanium, preferably 4.6 wt. % or less titanium, more preferably 4.2 wt. % or less titanium.
 15. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 3.2 wt. % or more tungsten.
 16. (canceled)
 17. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 1.2 wt. % or more molybdenum, preferably 1.7 wt. % or more molybdenum.
 18. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 6.1 wt. % or more cobalt, preferably 7.1 wt. % or more cobalt.
 19. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 24.6 wt. % or less cobalt, preferably 22.8 wt. % or less cobalt, more preferably 19.1 wt. % or even 16.8 wt % or 16.6 wt. % or less cobalt, even more preferably 14.7 wt. % or less cobalt, most preferably 11.1 wt. % or less cobalt.
 20. The nickel-based alloy composition of claim 1, consisting of, in weight percent, 11.9 wt. % or less chromium, preferably 11.5 wt. % or less chromium, more preferably 11.2 wt. % or less chromium.
 21. The nickel-based alloy composition of claim 1, consisting of, 51-62% volume fraction gamma prime at 850° C., preferably 51-56% volume fraction gamma prime at 850° C.
 22. A powder of the nickel-based alloy composition of claim
 1. 23.-25. (canceled) 